Welding Additive, Use of the Welding Additive and Component

ABSTRACT

A welding additive is provided. A component including a welding additive is also provided. The welding additive improves the weldability of a few nickel-based superalloys and includes the following contents (in wt %): 10.0%-20.0% chromium, 5.0%-15.0% cobalt, 0.0%-10.0% molybdenum, 0.5-3.5% tantalum, 0.0%-5.0% titanium, 1.5%-5.0% aluminum, 0.3%-0.6% boron, remainder nickel.

CROSS REFERENCE TO RELATED APPLICATIONS

This application is the US National Stage of International Application No. PCT/EP2009/061727, filed Sep. 10, 2009 and claims the benefit thereof. The International Application claims the benefits of European Patent Office application No. 08019282.6 EP filed Nov. 4, 2008. All of the applications are incorporated by reference herein in their entirety.

FIELD OF INVENTION

The invention relates to a weld filler as described in the claims, to the use thereof as claimed in the claims and to a component as claimed in the claims.

BACKGROUND OF INVENTION

Of all high-temperature materials, nickel-based superalloys have the most favorable combination of mechanical properties, resistance to corrosion and processability for gas turbine construction for aircraft and power plants. The considerable increase in strength is made possible in particular by the particle hardening with very high proportions by volume of the coherent γ′ phase Ni₃(Al—Ti, Ta, Nb). However, in general alloys with a higher γ′ content can only be considered weldable to a limited extent. This poor weldability is caused by:

a) Nickel alloys generally have a relatively low thermal conductivity and a relatively high coefficient of thermal expansion, similar to the values of austenitic steels and Co alloys. The welding heat which is introduced is therefore dissipated comparatively slowly, and the inhomogeneous heating leads to high thermal stresses, causing thermal fatigue which can only be dealt with at considerable effort. b) Nickel alloys are very sensitive to hot cracks in the event of a rapid change in the temperature cycles within the high temperature range. The cause is grain boundary fusion resulting from fluctuations in the chemical composition (segregations) or the formation of low-melting phases, such as sulfides or borides. c) Nickel alloys generally have a high proportion of the γ′ phase in a γ matrix. In the case of nickel-based superalloys for turbine components, the γ′ phase amounts to greater than 40 vol %. This achieves a high strength but also leads to a low ductility of the material, in particular at low temperatures and in the range of the temperature field in which the γ/γ′ precipitation phenomenon may occur (“ductility-dip temperature range”, also known as the “subsolidus ductility dip”, approximately 700° C. to 1100° C., depending on the alloy). Consequently, stresses which occur can less readily be absorbed through plastic flow, which generally increases the risk of crack formation. d) Nickel alloys exhibit the phenomenon of post-weld heat treatment cracks, also known as strain-age cracking. In this case, cracks are produced in a characteristic way in the first heat treatment following the weld as a result of γ/γ′ precipitation phenomena in the heat-affected zone or—if the weld filler can form the γ′ phase—also in the weld metal. This is caused by local stresses which form during the precipitation of the γ′ phase as a result of the contraction of the surrounding matrix. The susceptibility to strain-age cracking increases with an increasing level of γ′-forming alloy constituents, such as Al and Ti, since this also increases the proportion of γ′ phase in the microstructure.

If welds in which the base metal and the filler are identical are attempted at room temperature using conventional welding processes, for many industrial Ni-based superalloys for turbine laser vanes (e.g. IN 738 LC, Rene 80, IN 939), it is not currently possible to avoid the formation of cracks in the heat-affected zone and in the weld metal.

At present, a number of processes and process steps are known to improve the weldability of nickel-based superalloys:

a) Welding with Preheating: One way of avoiding cracks when welding nickel-based superalloys using high-strength fillers (likewise nickel-based superalloys) is to reduce the temperature difference and therefore the stress gradient between weld joint and the remainder of the component. This is achieved by preheating the component during the welding. One example is manual TIG welding in a shield and gas box, with the weld joint being preheated inductively (by means of induction coils) to temperatures of greater than 900° C. However, this makes the welding process significantly more complicated and expensive. Moreover, on account of inaccessibility, this cannot be implemented for all regions which are to be welded. b) Welding with Extremely Little Introduction of Heat: This involves the use of welding processes which ensure that very little heat is introduced into the base metal. These processes include laser welding and electron beam welding. Both processes are very expensive. Moreover, they require outlay on programming and automation, which may be uneconomical for repair welds, with frequently fluctuating damage patterns and locations.

US 2004/0115086 A1 has disclosed a nickel alloy with various additions.

SUMMARY OF INVENTION

Therefore, it is an object of the invention to provide a weld filler, a use of the weld filler, a welding process and a component which overcome the problems of the prior art.

The object is achieved by a weld filler as claimed in the claims, by the use of the weld filler as claimed in the claims and a component as claimed in the claims.

The subclaims give advantageous configurations which can advantageously be combined with one another as desired.

The invention proposes a weld filler and a use thereof which allows the repair welding of gas turbine blades or vanes and other hot-gas components made from nickel-based superalloys by manual or automated welding at room temperature. The weld filler is likewise a γ′-hardened nickel-based superalloy, but differs in particular from the material of a substrate of a component that is to be prepared. The welding repair allows a low cycle fatigue (LCF) corresponding to approximately 50% or more of the properties of the base metal (the weld withstands 50% of the LCF cycles of the base metal).

BRIEF DESCRIPTION OF THE DRAWINGS

The invention is explained in more detail below. In the drawing:

FIG. 1 shows a list of the composition of materials which can be welded using the filler according to the invention,

FIG. 2 shows a gas turbine,

FIG. 3 shows a perspective view of a turbine blade or vane, and

FIG. 4 shows a perspective view of a combustion chamber element.

DETAILED DESCRIPTION OF INVENTION

The invention proposes an increase in the boron fraction as an alloy element. The fraction of boron in the alloy should lie between 0.3 wt % and 0.6 wt %. Said increase is intended to improve the hot cracking resistance with fillers of the same and similar type. The filler, which allows repair welding of gas turbine blades and other hot-gas components made from Ni-based superalloys by manual or automated welding at room temperature.

Mechanism for Improving Resistance to Hot Cracking:

Hot cracks occur if, in the temperature range, the so-called BTR (Brittleness Temperature Range), the local deformability is too low to absorb the expansion caused by the welding. In relation to conventional Ni-based superalloys in which the boron fraction is limited to a maximum of 0.03 wt %, the boron eutectic completely encompasses the grain boundaries and, during solidification, acts as a type of “damper” or “buffer” and absorbs the deformation forces generated. Here, the melting temperature of the eutectic should be lower than the temperature Tu—temperature beyond which the deformation capability P increases again.

The base materials should not be overheated. Therefore, and use the welding parameters with low current intensity and fillers with small diameters are recommended.

Furthermore, the increased fraction of boron in the alloy increases the creep rupture strength and heat resistance with a parallel increase in resistance to oxidation in the aggressive media.

The invention proposes a welding process for welding components such as hot-gas components 138, 155 (FIG. 3, 4) and turbine blades or vanes 120, 130 (FIG. 2) made from nickel-based superalloys, which preferably includes the following characteristics:

-   -   Heat treatment prior to the welding with a view to coarsening γ′         phase in the base metal made from nickel-based superalloy (cf.         EP 1 428 897 A1). This heat treatment, also known as overageing,         increases the ductility and therefore the weldability of the         base metal.     -   Welding without preheating (at room temperature) using         conventional manual welding processes, such as TIG or plasma         powder welding, or alternatively welding using automated         processes, such as laser powder welding or automated plasma         powder welding, likewise at room temperature.     -   Use of closed shielding gas or vacuum boxes, into which the         entire component is introduced during welding, in order to         protect it from oxidation, is not required. There is also no         need for through-flow boxes, in which the component is protected         during welding by a correspondingly large flow of shielding gas.     -   For base metals which are extremely prone to hot cracking and/or         oxidation during welding, it is recommended to using shielding         gas which contains nitrogen to suppress the hot cracking and/or         hydrogen to reduce the oxidation (see EP 04011321.9).     -   Heat treatment after welding to homogenize base metal and weld         filler: solution annealing. The solution annealing temperature         should be adapted to the base metal. The solution annealing         temperature must be higher than the solution annealing         temperature but lower than the solidus temperature of the weld         filler. The single-stage or multi-stage age hardening to set the         desired γ′ morphology (size, shape, distribution) can take place         immediately afterwards or at a later stage during the processing         of the hot-gas components.

This weld filler has relatively good welding properties at room temperature. To achieve this, the levels of Al and Ti in the alloy were selected in such a way as to achieve a very low susceptibility to strain-age cracking. The chromium content is selected such that the alloy forms a corrosion-resistant Cr₂O₃ covering layer and contains a sufficient reservoir for regeneration of this layer under operating conditions.

Iron: Iron is preferably limited to at most 1.5 wt %, in order to improve the resistance of the alloy to oxidation and to reduce the risk of embrittling TCP phases (TCP=topologically closed packed) being formed.

Silicon: Silicon is preferably limited to at most 0.5 wt %, in order to minimize hot cracking.

Moreover, limits were determined for the optional constituents C, Fe, Mn, S, P, Hf, La, Si or Zr, in which an optimum between negative and positive influence is given.

When producing the component and during welding, oxides and in particular sulfides may form at the grain boundaries. These thin, intercrystalline eutectics containing sulfur and oxygen on the one hand embrittle the grain boundaries. On the other hand, they have a low melting temperature, which leads to a high susceptibility to grain boundary cracking as a result of local fusion of the grain boundaries.

The oxygen embrittlement is counteracted in particular by a local change in the chemical composition of the grain boundaries brought about by the addition of Hf, which segregates at the grain boundary and thereby makes grain boundary diffusion on the part of the oxygen more difficult, thus impeding grain boundary embrittlement, which is caused by oxygen. Moreover, hafnium is incorporated in the γ′ phase, increasing its strength.

The following table represents the invention (details in wt %).

Element Effect Cr 10.0-20.0 Corrosion resistance, increases the resistance to sulfidation, solid solution hardening Co  5.0-15.0 Reduces the stacking fault energy, resulting in increased creep strength, improves the solution annealing properties Mo  0.0-10.0 Solid solution hardening, increases the modulus of elasticity, reduces the diffusion coefficient Ta 0.0-3.5 Ti 0.0-5.0 Substitutes Al in γ′, increases the γ′ volume proportion Al 1.5-5.0 γ′ formation, only effective long-term protection against oxidation at > approx. 950° C., strong solid solution hardening Fe max 1.5 Promotes the formation of TCP phases, has an adverse effect on resistance to oxidation Mn max 0.1 Si max 0.5 Promotes the formation of TCP phases, increases hot cracking C  max 0.25 Carbide formation B 0.3-0.6 Element with grain boundary activity (large atom), increases the grain boundary cohesion, reduces the risk of incipient cracking, increases the ductility and creep rupture strength, prevents the formation of carbide films on grain boundaries, reduces the risk of oxidation Zr 0.0-0.1 Bonds S and C, increases resistance to hot cracking Hf 0.0-0.5 Reduces hot cracking capability during casting, is incorporated in γ′, increases the strength thereof, improves resistance to oxidation La 0.0-0.1 Bonds S, increases hot cracking capability S  max 0.015 Metallurgical impurity, increases hot cracking P  max 0.03 Metallurgical impurity, has an adverse effect on weldability Ni Remainder

One application example is the welding of the alloy Rene80, in particular when subject to operational stresses, by means of manual TIG welding and plasma-arc powder surfacing. Further welding processes and repair applications are not ruled out. The weld repair joints have properties which allow “structural” repairs in the airfoil/platform transition radius or in the airfoil of a turbine blade or vane.

Other nickel-based fillers can be selected according to the level of the γ′ phase, specifically for preference greater than or equal to 35 vol %, with a preferred maximum upper limit of 75 vol %.

The materials IN 738, IN 738 LC, IN 939, PWA 1483 SX or IN 6203 DS can preferably be welded using the weld filler according to the invention.

FIG. 2 shows, by way of example, a partial longitudinal section through a gas turbine 100.

In the interior, the gas turbine 100 has a rotor 103 with a shaft 101 which is mounted such that it can rotate about an axis of rotation 102 and is also referred to as the turbine rotor.

An intake housing 104, a compressor 105, a, for example, toroidal combustion chamber 110, in particular an annular combustion chamber, with a plurality of coaxially arranged burners 107, a turbine 108 and the exhaust-gas housing 109 follow one another along the rotor 103.

The annular combustion chamber 110 is in communication with a, for example, annular hot-gas passage 111, where, by way of example, four successive turbine stages 112 form the turbine 108.

Each turbine stage 112 is formed, for example, from two blade or vane rings. As seen in the direction of flow of a working medium 113, in the hot-gas passage 111 a row of guide vanes 115 is followed by a row 125 formed from rotor blades 120.

The guide vanes 130 are secured to an inner housing 138 of a stator 143, whereas the rotor blades 120 of a row 125 are fitted to the rotor 103 for example by means of a turbine disk 133.

A generator (not shown) is coupled to the rotor 103.

While the gas turbine 100 is operating, the compressor 105 sucks in air 135 through the intake housing 104 and compresses it. The compressed air provided at the turbine-side end of the compressor 105 is passed to the burners 107, where it is mixed with a fuel. The mix is then burned in the combustion chamber 110, forming the working medium 113. From there, the working medium 113 flows along the hot-gas passage 111 past the guide vanes 130 and the rotor blades 120. The working medium 113 is expanded at the rotor blades 120, transferring its momentum, so that the rotor blades 120 drive the rotor 103 and the latter in turn drives the generator coupled to it.

While the gas turbine 100 is operating, the components which are exposed to the hot working medium 113 are subject to thermal stresses. The guide vanes 130 and rotor blades 120 of the first turbine stage 112, as seen in the direction of flow of the working medium 113, together with the heat shield elements which line the annular combustion chamber 110, are subject to the highest thermal stresses.

To be able to withstand the temperatures which prevail there, they have to be cooled by means of a coolant.

Substrates of the components may likewise have a directional structure, i.e. they are in single-crystal form (SX structure) or have only longitudinally oriented grains (DS structure).

By way of example, iron-based, nickel-based or cobalt-based superalloys are used as material for the components, in particular for the turbine blade or vane 120, 130 and components of the combustion chamber 110.

Superalloys of this type are known, for example, from EP 1 204 776 B1, EP 1 306 454, EP 1 319 729 A1, WO 99/67435 or WO 00/44949.

The guide vane 130 has a guide vane root (not shown here), which faces the inner housing 138 of the turbine 108, and a guide vane head which is at the opposite end from the guide vane root. The guide vane head faces the rotor 103 and is fixed to a securing ring 140 of the stator 143.

FIG. 3 shows a perspective view of a rotor blade 120 or guide vane 130 of a turbo machine, which extends along a longitudinal axis 121.

The turbo machine may be a gas turbine of an aircraft or of a power plant for generating electricity, a steam turbine or a compressor.

The blade or vane 120, 130 has, in succession along the longitudinal axis 121, a securing region 400, an adjoining blade or vane platform 403 and a main blade or vane part 406 and a blade or vane tip 415.

As a guide vane 130, the vane 130 may have a further platform (not shown) at its vane tip 415.

A blade or vane root 183, which is used to secure the rotor blades 120, 130 to a shaft or a disk (not shown), is formed in the securing region 400.

The blade or vane root 183 is designed, for example, in hammerhead for Other configurations, such as a fir-tree or dovetail root, are possible.

The blade or vane 120, 130 has a leading edge 409 and a trailing edge 412 for a medium which flows past the main blade or vane part 406.

In the case of conventional blades or vanes 120, 130, by way of example solid metallic materials, in particular superalloys, are used in all regions 400, 403, 406 of the blade or vane 120, 130.

Superalloys of this type are known, for example, from EP 1 204 776 B1, EP 1 306 454, EP 1 319 729 A1, WO 99/67435 or WO 00/44949.

The blade or vane 120, 130 may in this case be produced by a casting process, also by means of directional solidification, by a forging process, by a milling process or combinations thereof.

Work pieces with a single-crystal structure or structures are used as components for machines which, in operation, are exposed to high mechanical, thermal and/or chemical stresses.

Single-crystal work pieces of this type are produced, for example, by directional solidification from the melt. This involves casting processes in which the liquid metallic alloy solidifies to form the single-crystal structure, i.e. the single-crystal work piece, or solidifies directionally.

In this case, dendritic crystals are oriented along the direction of heat flow and form either a columnar crystalline grain structure (i.e. grains which run over the entire length of the work piece and are referred to here, in accordance with the language customarily used, as directionally solidified) or a single-crystal structure, i.e. the entire work piece consists of one single crystal. In these processes, a transition to globular (polycrystalline) solidification needs to be avoided, since non-directional growth inevitably forms transverse and longitudinal grain boundaries, which negate the favorable properties of the directionally solidified or single-crystal component.

Where the text refers in general terms to directionally solidified microstructures, this is to be understood as meaning both single crystals, which do not have any grain boundaries or at most have small-angle grain boundaries, and columnar crystal structures, which do have grain boundaries running in the longitudinal direction but do not have any transverse grain boundaries. This second form of crystalline structures is also described as directionally solidified microstructures (directionally solidified structures).

Processes of this type are known from U.S. Pat. No. 6,024,792 and EP 0 892 090 A1.

The blades or vanes 120, 130 may likewise have coatings protecting against corrosion or oxidation, for example (MCrAlX; M is at least one element selected from the group consisting of iron (Fe), cobalt (Co), nickel (Ni), X is an active element and represents yttrium (Y) and/or silicon and/or at least one rare earth element, or hafnium (Hf)). Alloys of this type are known from EP 0 486 489 B1, EP 0 786 017 B1, EP 0 412 397 B1 or EP 1 306 454 A1.

The density is preferably 95% of the theoretical density.

A protective aluminum oxide layer (TGO=thermally grown oxide layer) forms on the MCrAlX layer (as intermediate layer or as outermost layer).

It is also possible for a thermal barrier coating, which is preferably the outermost layer and consists for example of ZrO₂, Y₂O₃—ZrO₂, i.e. unstabilized, partially stabilized or fully stabilized by yttrium oxide and/or calcium oxide and/or magnesium oxide, to be present on the MCrAlX.

The thermal barrier coating covers the entire MCrAlX layer.

Columnar grains are produced in the thermal barrier coating by means of suitable coating processes, such as for example electron beam physical vapor deposition (EB-PVD).

Other coating processes are conceivable, for example atmospheric plasma spraying (APS), LPPS, VPS or CVD. The thermal barrier coating may have porous, microcrack-containing or macrocrack-containing grains for better thermal shock resistance. The thermal barrier coating is therefore preferably more porous than the MCrAlX layer.

The blade or vane 120, 130 may be hollow or solid in form. If the blade or vane 120, 130 is to be cooled, it is hollow and may also have film-cooling holes 418 (indicated by dashed lines).

FIG. 4 shows a combustion chamber 110 of the gas turbine 100. The combustion chamber 110 is configured, for example, as what is known as an annular combustion chamber, in which a multiplicity of burners 107 arranged circumferentially around the axis of rotation 102 open out into a common combustion chamber space 154 and which generate flames 156. For this purpose, the combustion chamber 110 overall is of annular configuration positioned around the axis of rotation 102.

To achieve a relatively high efficiency, the combustion chamber 110 is designed for a relatively high temperature of the working medium M of approximately 1000° C. to 1600° C. To allow a relatively long service life even with these operating parameters, which are unfavorable for the materials, the combustion chamber wall 153 is provided, on its side which faces the working medium M, with an inner lining formed from heat shield elements 155.

On account of the high temperatures in the interior of the combustion chamber 110, it is also possible for a cooling system to be provided for the heat shield elements 155 and/or for their holding elements. The heat shield elements 155 are in this case for example hollow and may also have cooling holes (not shown) opening out into the combustion chamber space 154.

On the working medium side, each heat shield element 155 is equipped with a particularly heat-resistance protective layer (McrAlX layer and/or ceramic coating) or is made from material that is able to withstand high temperatures (solid ceramic bricks).

These protective layers may be similar to the turbine blades or vanes, i.e. for example McrAlX: M is at least one element selected from the group consisting of iron (Fe), cobalt (Co), nickel (Ni), X is an active element and represents yttrium (Y) and/or silicon and/or at least one rare earth element, or hafnium (Hf). Alloys of this type are known from EP 0 486 489 B1, EP 0 786 017 B1, EP 0 412 397 B1 or EP 1 306 454 A1.

It is also possible, for example, for a ceramic thermal barrier coating to be present on the McrAlX, consisting for example of ZrO₂, Y₂O₃—ZrO₂, i.e. unstabilized, partially stabilized or fully stabilized by yttrium oxide and/or calcium oxide and/or magnesium oxide.

Columnar grains are produced in the thermal barrier coating by means of suitable coating processes, such as for example electron beam physical vapor deposition (EB-PVD).

Other coating processes are conceivable, for example atmospheric plasma spraying (APS), LPPS, VPS or CVD. The thermal barrier coating may have porous, microcrack-containing or macrocrack-containing grains for better thermal shock resistance.

Refurbishment means that after they have been used, protective layers may have to be removed from turbine blades or vanes 120, 130, heat shield elements 155 (e.g. by sand-blasting). Then, the corrosion and/or oxidation layers and products are removed.

If appropriate, cracks in the turbine blade or vane 120, 130 or the heat shield element 155 are also repaired using the weld filler according to the invention. This is followed by recoating of the turbine blades or vanes 120, 130, heat shield elements 155, after which the turbine blades or vanes 120, 130 or the heat shield elements 155 can be reused. 

1.-18. (canceled)
 19. A weld filler, comprising (in wt %): 10.0%-20.0% chromium; 5.0%-15.0% cobalt; 0.0%-10.0% molybdenum; 0.5%-3.5% tantalum; 0.0%-5.0% titanium; 1.5%-5.0% aluminum; and 0.3%-0.6% boron;
 20. The weld filler as claimed in claim 19, further comprising at least 0.01 wt % carbon.
 21. The weld filler as claimed in claim 19, further comprising at least 0.1 wt % iron.
 22. The weld filler as claimed in claim 19, further comprising at least 0.01 wt % manganese.
 23. The weld filler as claimed in claim 19, further comprising at least 0.005 wt % sulfur.
 24. The weld filler as claimed in claim 19, further comprising at least 0.005 wt % phosphorus.
 25. The weld filler as claimed in claim 19, further comprising at least 0.05 wt % hafnium.
 26. The weld filler as claimed in claim 19, further comprising at least 0.01 wt % lanthanum.
 27. The weld filler as claimed in claim 19, further comprising at least 0.05 wt % zirconium.
 28. The weld filler as claimed in claim 19, further comprising at least 0.05 wt % silicon.
 29. The weld filler as claimed in claim 19, consisting of nickel, chromium, cobalt, molybdenum, titanium, tantalum, aluminum and boron.
 30. The weld filler as claimed in claim 19, consisting of nickel, chromium, cobalt, molybdenum, titanium, tantalum, aluminum, boron, and at least two elements selected from the group consisting of carbon, iron, manganese, silicon, zirconium, hafnium and lanthanum.
 31. The weld filler as claimed in claim 19, consisting of nickel, chromium, cobalt, molybdenum, titanium, tantalum, aluminum, boron, and at least three elements selected from the group consisting of carbon, iron, manganese, silicon, zirconium, hafnium and lanthanum.
 32. The weld filler as claimed in claim 19, consisting of nickel, chromium, cobalt, molybdenum, titanium, tantalum, aluminum, boron, carbon, iron, manganese, silicon, zirconium, hafnium and lanthanum.
 33. A nickel-based component including a weld filler as claimed in claim 19, wherein a nickel-based material is different than the weld filler.
 34. The component as claimed in claim 33, wherein the nickel-based material includes a γ′-phase in a proportion of ≧40 vol %.
 35. The component as claimed in claim 33, wherein the nickel-based material of the component includes IN 738 or IN 738 LC, Rene 80 or IN939 and also PWA 1483 SX or IN 6203 DS.
 36. A weld filler, comprising (in wt %): 10.0%-20.0% chromium; 5.0%-15.0% cobalt; 0.0%-10.0% molybdenum; 0.5%-3.5% tantalum; 0.0%-5.0% titanium; 1.5%-5.0% aluminum; and 0.3%-0.6% boron; at most 0.30% carbon; at most 1.8% iron; at most 0.15% manganese; at most 0.03% sulfur; at most 0.06% phosphorus; at most 0.7% hafnium; at most 0.2% lanthanum; at most 0.7% silicon; at most 0.2% zirconium; and remainder nickel.
 37. The weld filler as claimed in claim 36, further comprising at least 0.05 wt % carbon.
 38. The weld filler as claimed in claim 36, further comprising at least 0.35 wt % iron. 